University of Cincinnati

UNIVERSITY OF CINCINNATI

Date:______

I, ______, hereby submit this work as part of the requirements for the degree of: in:

It is entitled:

This work and its defense approved by:

Chair: ______

EFFECT OF AGING ON ABRASIVE WEAR RESISTANCE OF SILICON CARBIDE

PARTICULATE REINFORCED ALUMINUM MATRIX COMPOSITE

A thesis submitted to the

Division of Graduate Studies and Research of the

University of Cincinnati

In partial fulfillment of the requirement for the degree of

MASTER OF SCIENCE

in the Department of Chemical and Materials Engineering

of the College of Engineering

2007

Varun Sethi

B.Tech, National Institute of Technology, Jamshedpur, 2005

Committee Chair: Dr. R Y. Lin

Abstract

The effect of aging on the wear resistance of SiC particle reinforced aluminum composites was investigated. The as cast Al/SiCp composite used in this study was purchased from Duralcan with A-356 matrix and 23Vol% of SiC reinforcement. This composite was solutionized at 565ºC and then aged at 180 ºC for different time intervals and changes in hardness and wear resistance were measured using a Rockwell B hardness tester and Pin-on-disc wear tester respectively. For reference purpose an alloy of Al-10wt% silicon was used. Wear resistance of the aged composites was found superior to the as cast composite with the peak aged composite showing the maximum wear resistance and overaging resulted in a decrease in wear resistance. Results also showed that the wear resistance of the composites was greater than the monolithic alloy at all loads and wear rate was found to increase with pressure. The wear resistance of

Al-Si(10wt%) alloy was found to increase with aging, but no variation in wear rate was found among the aged alloys. Scanning electron micrographs of worn surfaces of composites revealed that the principal mechanism of matrix removal was microcutting and microcracking. Very few SiC reinforcements were found on the worn surfaces suggesting that that the penetration depth of the abrasive was greater than the reinforcement particle size in most cases. Finally the differences in wear resistance of the composites were rationalized on the basis of changes at the interface of SiC particle and aluminum matrix due to the presence of precipitates.

Acknowledgement

I would like to take this opportunity to express my sincere appreciation for my research advisor and thesis committee chair, Dr. Ray Y. Lin, for his constant support, encouragement, and able guidance throughout this study.

I am grateful to my thesis committee members, Dr. Jude Iroh, and Dr. Rodney Roseman for their review and helpful criticism.

I would like to thank Mr. Ratandeep Kukreja and Dr. Doug Kohls for their help in electron microscopy. I would like to appreciate CME staff Dale and Molly for their help. I am also grateful to the CME department for extending all of its resources.

Table of Contents

Page

List of Tables i

List of Figures ii

List of symbols iv

1.0 Introduction 1

2.0 Background 3

2.1 Composite materials 3

2.2 Metal Matrix Composites 5

2.3 Aluminum Based Composites 6

2.4 Wear applications of aluminum based

MMC’s in automotive industry 7

2.5 Fabrication of Aluminum matrix composites 11

2.5.1 Solid State Processing 11

2.5.2 Liquid State Processing 12

2.6 Interfacial Reaction 14

2.7 A356 Casting alloy 18

2.8 Aging Behavior of Metal Matrix Composites 21

2.9 Abrasive Wear 24

2.9.1 Hardness 25

2.9.2 Effect of abrasive grit dimension 26

2.9.3 Fracture toughness 27 2.10 Wear Resistance of Metal Matrix Composites 28

2.10.1 Nature of the reinforcement 30

2.10.2 Effect of reinforcement 32

2.10.3 Effect of increasing volume fraction of

reinforcement on wear 34

2.10.4 Interfacial strength 36

2.10.5 Shape of dispersoid and reinforcement 37

2.10.6 Particle size 38

3.0 Experimental 42

3.1 Materials 42

3.2 Chemical analysis of composite 42

3.3 Particle size measurement 43

3.4 Density Measurement 44

3.5 Wear Test 45

3.6 Heat treatment 47

3.7 Hardness Testing 47

3.8 X-ray Diffraction 47

3.9 Microscopic Examination 48

4.0 Results 49

4.1 Composite 49

4.2 Aging Studies 51

4.3 X-ray Diffraction 58

4.4 EDS analysis of precipitate 59 4.5 Wear Behavior 60

4.6 Worn Surfaces 61

5.0 Discussion 67

5.1 Aging Studies 70

5.2 X-ray Diffraction analysis 73

5.3 Wear Studies 74

5.3.1 Effect of load on wear 74

5.3.2 Mechanism of matrix removal 75

5.3.3 Mechanism of particle removal 76

5.3.4 Dominant Mechanism 81

5.3.5 Comparison of wear behavior of

composite with alloy 87

5.3.6 Comparison with previous studies 87

6.0 Conclusions 90

7.0 Future work 85

8.0 References 86

List of Tables

Page

Table 1 -- Selected Cast Composite Components with proven

automobile applications 9

Table 2 -- Properties of A356 casting alloy 19

Table 3 -- Nominal composition of A356 aluminum alloy 42

Table 4 -- Hardness of the composites and alloys 52

Table 5 -- EDS result of precipitate present in the composite

after aging 60

List of Figures

Page

Figure 2.1 -- Aluminum-Silicon phase diagram 20

Figure 2.2 -- Variation of matrix microhardness as a function of aging

time at 177°C 23

Figure 2.3 -- Specific wear rate in several aluminum base particulate

composites sliding against steel as a function of particle

volume fraction 31

Figure 2.4 -- Drawing illustrating the concept of dimensions of contact area. 39

Figure 3.1 -- Schemetic representation of pin-on-disc wear tester 46

Figure 4.1 -- Microstructure of the composite showing (a) uniform

distribution of SiC particles in the matrix and

(b) SiC particle embedded in aluminum matrix 49

Figure 4.2 -- (a) Optical Micrograph of the composite at 1500X

(b) Optical Micrograph of the scale micrometer at 1500X 50

Figure 4.3 -- Variation of Rockwell hardness with aging time 53

Figure 4.4 -- SEM Micrographs of the composites after aging at 180ºC 54

Figure 4.5 -- SEM Micrographs of the alloys after aging at 180ºC 56

Figure 4.6 -- X-ray diffraction pattern of as cast and aged composite 58

Figure 4.7 -- SEM Micrograph of a precipitate for EDS analysis 59

Figure 4.8 -- Variation of Wear Rate with Pressure for composites 62

Figure 4.9 -- Variation of Wear Rate with Pressure taking the average of three 63

tests at a sliding distance of 80m 58

Figure 4.10 -- Variation of Wear Rate with Pressure for alloys 64

Figure 4.11 -- SEM micrograph of the 12 hours aged composite surface after wear 65

Figure 4.12 -- SEM micrographs of composites and alloy showing

microcutting phenomenon 66

Figure 4.13 -- SEM micrographs of the composites and alloy showing

Microcracking 68

Figure 5.1 -- Figure 5.1 Free Energy vs. Composition diagram for α and β

phases in Al-Si System 72

Figure 5.2 -- SEM Micrograph showing the fracture of the particle 77

Figure 5.3 -- SEM micrograph of the composite showing weakening

of the interfacial bond 78

Figure 5.4 -- SEM Micrograph of the composite showing a pull-out of a

SiC particle from the Matrix 78

Figure 5.5 -- SEM micrographs of Aged composites showing adhesion of

precipitate with SiC Particle 80

Figure 5.6 -- Schemetic showing adherence of the precipitates with the reinforcement 82

Figure 5.7 -- Schematic diagrams showing the interaction of the abrasive with

the SiC particle 84

Figure 5.8 -- Schematic diagrams showing interactions between an abrasive

particle and a SiC, composite at penetration depths of

(a) h < d and (b) h> d. 86

iii

Figure 5.9 -- The relative wear resistance vs. relative abrasive penetration depth. 86

Figure 5.10 -- SEM Micrograph showing a protruding SiC particle 88

List of Symbols

ρ Density

Vr Volume fraction of reinforcement

Vm Volume fraction of matrix

P Applied normal load

H Bulk hardness of material

Heff Effective hardness

Hm Hardness of the work hardened matrix phase

Hs Hardness of the second phase particles

Wc Wear rate of composite

Wp Wear rate of particle

Wal Wear rate of aluminum matrix fal Area fraction of aluminum phase

fap Area fraction of particle phase

S Sliding distance

d Particle size

Vf Volume fraction of the particle

1.0 Introduction

Metal matrix composites are materials with metals as the base and distinct, typically ceramic phases added as reinforcements to improve the properties. The reinforcements can be in the form of fibers, whiskers and particulates. Properties of the metal matrix composites can be tailored by varying the nature of constituents and their volume fraction. They offer superior combination of properties in such a manner that today no existing monolithic material can rival.

They are increasingly being used in the aerospace and automobile industries because of their improved strength, stiffness and increased wear resistance over unreinforced alloys.

Aluminum is the most popular matrix for metal matrix composites because of its low density, its capability to be strengthened by precipitation, good corrosion resistance, high thermal and electrical conductivity, improved tribological properties over monolithic alloys and its high damping capacity. They are usually reinforced with SiC, Al2O3 and carbon. They find applications in automobile, defense and aerospace sectors because of their excellent combination of higher specific strength and stiffness, improved wear and seizure resistance, higher elevated temperature strength over their base alloys. Aluminum matrix composites find potential applications in automobile components like piston, cylinder liner, brake drums, crankshafts, etc[1-5]. The key benefits of aluminum matrix composites in the transportation sector are lower fuel consumption, economic and environmental benefits. The above mentioned components components undergo sliding as well as abrasive type of wear against the counter surface during operation. In view of this considerable attentions have been paid on assessing the sliding [6-11] and abrasive wear [12-17] behavior of aluminum matrix composites under varying tribological conditions. Most of these studies have been focused on the influence of the type, volume

fraction, size and the geometry of the reinforcing elements on wear behavior. Very little attention has been directed towards the effect of ageing on the wear resistance. Lin and Liu [18] reported that wear resistance of the SiC particulate reinforced Al-Zn-Mg alloy can be altered by ageing, with the highest wear resistance observed after overaging. These authors attributed this to the increase in particle/matrix interfacial strength with overaging. However this study was in conflict with the results proposed by Song [19] who found that the wear resistance decreases with overaging due to the coarsening of precipitates and loss in hardness.

In the present investigation an effort has been made to relate the changes in the microstructure caused by ageing to the changes in the wear behavior of composite. In order to understand this phenomenon, wear tests of the as-cast and aged-composites were carried out using a 60 grit abrasive counterface (268µm particle size) at different loads and a study of the composite microstructure after aging and wear tests was carried out using scanning electron microscopy.

2.0 Background

2.1 Composite Materials

In recent years there has been a rapid advance in the field of composite materials. A complete review of composite materials would have to cover large sections of material science and engineering, metallurgy, polymer technology, fracture mechanics and applied mechanics.

Extensive literature is available and hence only a brief introduction is given here.

The term „composite‟ in composite materials means that two or more materials are combined together in a certain order on a macroscopic level to form a new material with different and attractive properties. The reason alloys are not considered in this category is because they are homogeneous on the macroscopic level. Composite materials consist of a bulk material called the matrix, and a filler of some types such as fibers, whiskers or particles. The choice of materials for matrices and reinforcements is almost limitless. Nearly any conventional engineering material can be used as a matrix, while new reinforcements are announced almost monthly. The matrix is responsible for transferring the load to the reinforcement and for protecting the reinforcement surface from abrasion and environmental corrosion, both of which can initiate fracture. The matrix also determines much of the shear and composite properties, creep and flow properties, dielectric properties and thermomechanical properties of the composites. The reinforcement on the other hand is primarily responsible for the high specific strengths associated with composites. Composite materials are conventionally classified into three categories viz. polymer matrix, metal matrix and ceramic matrix depending on the matrix employed. Furthermore, composites can be grouped on the type of reinforcement provided [20].

They are:

Dispersion strengthened: In this type of composite, fine hard particulates, uniformly distributed, ranging in size from 0.01μm to 0.1μm and in a volume percentage from 1 to15% are used to increase the strength and hardness.

Particle reinforced: This is similar to the dispersion-strengthened category but the size of the particles is greater than 0.1μm and the volume percentages can be greater than 25%.

Fiber reinforced: This category includes all types of fibers, whiskers and filaments, continuous and non-continuous, over the whole range of concentration of the reinforcements. Although the continuous-fiber-reinforced metal matrix composites (MMCs) have the best combination of mechanical properties, particulate-reinforced MMCs are superior from the viewpoint of cost- performance trade-off. Compared to fiber reinforced MMCs, the particulate-reinforced MMCs possess improved ductility, reduced anisotropy of mechanical properties, as well as ease of secondary working with conventional metalworking techniques. The availability of inexpensive particulates as reinforcements serves as an added advantage.

The properties of a composite depend on the following

• Properties of constituent phases

• Relative amounts of constituents

• Geometry of the dispersed phase

• Shape of particles

• Particle size

• Particle distribution

• Interfacial reactions between constituents

Several properties of the composites may be assessed by the rule of mixture (ROM), which states that the property of the composite is the sum of products of the property of an individual constituent and its volume fraction in the composite. For example:

Density: ρc = ρrVr + ρmVm (1)

Where, ρ stands for density, V Stands for Volume fraction and subscripts c, r and m stand for composite, reinforcement and matrix respectively. This theoretical calculation can be used to check whether the composite produced has the optimum properties. Modulus, Strength,

Electrical properties may sometimes be estimated with ROM as the first approximation. Only particulate reinforced metal matrix composites will be dealt with here. Discussion on other types of composites is beyond the scope of this chapter.

2.2 Metal-Matrix Composites

Metal-matrix composites (MMCs) have, in the last three decades or so, come up in a big way primarily because of their superior mechanical properties compared to monolithic materials.

The principal advantage MMCs enjoy over other materials lies in the improved strength and hardness on a unit weight basis. Metal-matrix composites were first developed for the application in aerospace industries. The expansion into the non-aerospace and non-military areas came about as the price of MMCs came down largely due to the availability of inexpensive particulates and development of low cost fibers.

Metal-matrix composites have several advantages over conventional structural materials.

These advantages include a combination of the following properties [20]:

•High strength

•High modulus

•High toughness and impact properties

•Low sensitivity to temperature changes and thermal shock

•High surface durability and low sensitivity to surface flaws and

•High electrical and thermal conductivity.

Resistance to severe environments, toughness and retention of strength at elevated temperatures are some areas where metal-matrix composites perform better than other composites.

Metal-matrix composites are artificially produced by two or more materials having significantly different properties. The constituent materials that often do not posses very good properties on their own, when combined, produce some very attractive propositions for materials design. Metal-matrix composites find applications in the manufacture of aircraft, automobile engines, electrical machinery, electrical and thermalconductors, bearing materials, missiles, spacecrafts, rocket launch vehicle structures and numerous other applications.

2.3 Aluminum Based Composites

Aluminum is the most popular matrix for the metal matrix composites (MMCs). The Al alloys are quite attractive due to their low density, their capability to be strengthened by precipitation, their good corrosion resistance, high thermal and electrical conductivity, improved tribological properties over monolithic alloys and their high damping capacity. Aluminum matrix composites (AMCs) have been widely studied since the 1920s and are now used in sporting goods, electronic packaging, armours, aerospace and automotive industries. They offer a large variety of mechanical properties depending on the chemical composition of the Al-matrix. They

are usually reinforced by Al2O3, SiC, C but SiO2, B, BN, B4C, AlN may also be considered. The aluminum matrices are in general Al-Si,Al-Cu, 2xxx or 6xxx alloys. Aluminum composites has several improvements over monolithic aluminum, such as higher modulus, higher fatigue resistance, higher hardness, magnitudes of improvement in wear resistance, lower coefficient of thermal expansion, better high temperature strength, and greatly increased resistance to creep. In addition these composites still have the inherent useful properties of aluminum, such as superior corrosion resistance, low cost and good fracture toughness.

2.4 Wear applications of aluminum based MMC’s in automotive industry

Since the present study is based on wear testing of SiC reinforced Al-356, it becomes imperative to discuss the wear applications of this composite. In the automotive industry, the major driving forces for developing and implementing new materials and manufacturing technology are fuel economy, reduced vehicle emissions, and increased vehicle safety at competitive cost [21]. Light-weight materials such as Al–matrix composites permit lighter engine and structural components with improved properties and performance to be designed and used in automobiles. Replacing cast iron engine components with light-weight Al alloys requires overcoming the poor adhesion and seizure resistance of aluminum by dispersing SiC, Al2O3 or graphite particles in Aluminum. Considerable reduction in wear and friction is achieved by use of these particulates. Over the last decade, the enabling technologies to produce components from Al-based MMC materials have matured to a level where commercialization of products made from the new material has become feasible [22-31]. The auto industry has successfully applied Al-based particulate composites, chiefly SiC/Al and Al/ Al2O3, in pistons, engine blocks, disc rotor brakes, drums, calipers, connecting rods, internal combustion engine cylinder liners,

drive shafts, snow tire studs and other parts. Most notable example is the development of the all aluminum engine block by Honda [32] in the early 1990s. Here the Honda has produced a thin- walled cylindrical ceramic perform made of hybrid material of short alumina and carbon fibers, and squeeze infiltrated molten aluminum alloy to produce the engine block. The new engine block features is more compact with significant weight reductions compared with cast iron engine blocks and those made of aluminum alloy with cast iron liners, thus providing higher performance. Aluminum composites have superior thermal conductivity and lower density than cast iron, and this has been profitable in disc brake. According to Rusnak[33] during moderate to heavy braking, thermal conductivity plays a predominant role in determining peak rotor temperatures. Aluminum MMC rotors with their high thermal conductivity will run cooler and should show superior frictional stability over cast iron rotors. Aluminum composite brake rotors provide up to 60% weight reduction when compared to cast iron[34]. Under both sliding and abrasion wear, MMC displays lower wear rate which decreases as the reinforcement content increases. Depending upon the particle loading, abrasive wear rates are reduced 55–90% when compared to the wear rate of unreinforced Al, and at approximately 20 vol. %. SiC in Al, Al composite brake rotors have lower wear rate than cast iron [34]. Aluminum MMC brake discs are now in production with properties of both the disc material and friction lining material being refined to meet friction, wear and fabrication criteria[35,36].Aluminum based MMC‟s have also found applications in internal combustion engine cylinder liners. In this application hybrid composite containing both SiC and graphite particulates in Al-Si alloy matrices have proved to be successful candidate replacement materials, with advantageous low-load wear and overall seizure resistance [37]. Diesel engine pistons containing Saffil (Al2O3) short fibers have been in use by Toyota since 1985[21]. Reinforcing piston crown with a MMC reduces the piston

thickness at the crown and the overall piston weight. Additionally, the ceramic reinforcement reduces heat losses because of its high thermal resistance. Composite liners have better scuffing characteristics than conventional cast iron liners of the engine block. Table1 [21] lists the proven applications of some Aluminum based MMC‟s in the automotive industry

Table 1

Selected Cast Composite Components with proven automobile applications

______

Manufacturer Component and Composite

______

Duralcan, Martin Marietta, Pistons, Al/SiCp Lanxide

Duralcan, Lanxide Brake rotors, calipers, liners, Al/SiCp

GKN, Duralcan Propeller shaft, Al/SiCp

Nissan Connecting rod, Al/SiCw

Dow Chemical Sprockets, pulleys, covers,Mg/SiCp

Toyota Piston rings, Al/Al2O3 (saffil) & Al/Boriaw

Dupont, Chrysler Connecting rods, Al/Al2O3

Hitachi Current collectors, Cu/graphite

Associated Engineering, Inc. Cylinders, pistons, Al/graphite

Martin Marietta Pistons, connecting rods, Al/TiCp

Zollner Pistons, Al/fiberfrax

______

Manufacturer Component and Composite

______

Honda Engine blocks, Al/Al2O3 – Cf

Lotus Elise, Volkswagon Brake rotors, Al/SiCp

Chrysler Brake rotors, Al/SiCp

GM Rear brake drum for EV-1, driveshaft, engine cradle, Al/SiC

3M Missile fins, aircraft electrical access door, Al/Nextelf

Knorr-Bremse; Kobenhavn Brake disc on ICE bogies, SiC/Al

Alcoa Innometalx Multichip electronic module, Al/SiCp

Lanxide PCB Heat sinks, Al/SiCp

Cercast Electronic packages, Al/graphite foam

Textron Specialty Materials PCB heat sinks, Al/B ______

2.5 Fabrication of Aluminum Matrix Composites

Primary processes for manufacturing of aluminum matrix composites at industrial scale can be classified into two main groups [38].

1) Solid state processes

2) Liquid state processes

Powder blending followed by consolidation, diffusion bonding and vapor deposition techniques come under solid state processing. Liquid state processes include stir casting or compo casting, infiltration, spray casting and in situ (reactive) processing. The selection of the processing route depends on many factors including type and level of reinforcement loading and the degree of microstructural integrity desired.

2.5.1 Solid State Processing

Powder blending and consolidation (PM processing): This process involves the mixing of prealloyed aluminum powder with ceramic reinforcement particulates (or whiskers) in the desired volume ratio. Blending can be carried out in dry or in liquid suspension. Blending is usually followed by cold compaction, canning, degassing and high temperature consolidation stage such as hot isostatic pressing and extrusion. Powder metallurgy processed aluminum matrix composites contain oxide particles in the form of plate-like particles of few tens of nm thick and in volume fractions ranging from 0.05 to 0.5 depending on powder history and processing conditions. These fine oxide particles tend to act as a dispersion strengthening agent and often have a strong influence on the matrix properties particularly during heat treatment. The advantages in this process include the latitude in controlling the volume ratio of the constituents.

Diffusion Bonding: Mono filament- reinforced aluminum matrix composites are mainly produced by the diffusion bonding (foil- Fiber-foil) route or by the evaporation of relatively thick layers of aluminum on the surface of the fiber. 6061 Al-boron fiber composites have been produced by diffusion bonding via the foil-fiber foil process. However, the process is more commonly used to produce Ti based fiber reinforced composites. The process is cumbersome and obtaining high fiber volume fraction and homogenous fiber distribution is difficult. The process is not suitable to produce complex shapes and components.

Physical Vapor deposition: The process involves continuous passage of fiber through a region of high partial pressure of the metal to be deposited, where condensation takes place so as to produce a relatively thick coating on the fiber. The vapor is produced by directing a high power electron beam onto the end of a solid bar feed stock. Typical deposition rates are 5-10µm per minute. Composite fabrication is usually completed by assembling the coated fibers into a bundle or array and consolidating in a hot press operation. Composites with uniform distribution of fiber and volume fraction as high as 80% can be produced by this technique.

2.5.2 Liquid state processing

Stir Casting: This involves incorporation of ceramic particulate into liquid aluminum melt and allowing the mixture to solidify. The crucial thing in this process is to create good wetting between the particulate reinforcement and the liquid aluminum alloy melt. The simplest and most commercially used technique is known as vortex technique or stir casting technique which involves the introduction of pre-treated ceramic particles into the vortex of molten alloy created

by rotating impeller. It is possible to incorporate upto 30% ceramic particles in the size range 5-

100µm in a variety of molten aluminum alloys. Particle agglomeration, sedimentation in the melt and subsequently during solidification, and inhomogenity in reinforcement distribution due to the interaction between suspended ceramic particles and moving solid-liquid interface during solidification are some disadvantages in these cast composites.

Liquid Metal Infiltration: In this process liquid aluminum alloy is injected/ infiltrated into the interstices of porous performs of continuous fiber/ short fiber, whisker or particle to produce aluminum matrix composites. Depending on the nature of reinforcement and its volume fraction perform can be infiltrated, with or without the application of pressure or vaccum. Composites having reinforcement volume fraction ranging from 10 to 70% can be produced using a variety of infiltration techniques. The major advantages include low cost, easy processing steps and dense products.

Spray Deposition Processes: In these processes droplets of molten metal are sprayed together with the reinforcing phase and collected on a substrate where metal solidification is completed.

Alternately the reinforcement may be placed on the substrate, and molten metal may be sprayed onto it. The critical parameters in this process are the initial temperature, size distribution and velocity of metal drops, the velocity, temperature and feeding rate of the reinforcement (if it is simultaneously injected), and the position, nature , and the temperature of the substrate collecting the material. Advantages include fine grain size and low segregation in the matrix and minimization of interfacial reaction because the metal and the reinforcement contact only briefly.

In-Situ Processing( Reactive processing) : The term in-situ composites was first used for materials produced by solidification of polyphase alloys. When polyphase alloys solidify directionally with a plane front, they may exhibit a fine lamellar or rod like structure of β phase in an α phase matrix and the interphase spacing is the function of the growth rate. Some examples include XD process in which Ti, B and Al powders are heated to 800C to form TiB2 reinforced Aluminum. TiC –reinforced Al-Cu alloys have been obtained by bubbling CH4 and argon gas through a melt of Al-Cu-Ti. DIMOX process involves the directional oxidation of aluminum. In this process the alloy of Al-Mg is placed on the top of ceramic perform in a crucible. The entire assembly is heated to a suitable temperature in the atmosphere of free bearing gas mixture. Al-Mg alloy soon after melting infiltrates into the perform and composite is formed. A major advantage of this process is that the reinforcing phase is generally homogenously distributed, and spacing or size of the reinforcement may be adjusted in several cases by the solidification or reaction time. However, the choice of systems and the orientation of the reinforcement are limited, and kinetics of the processes, or the shape of the reinforcing phase, are sometimes difficult to control.

2.6 Interfacial Reaction

Aluminum/silicon carbide composites are non-equilibrium systems; consequently, chemical potential gradients exist at the Al/SiC interface. These gradients are a driving force for interfacial reactions that may occur during fabrication of the composites at high temperature or during heat treatment. The formation of aluminum carbide has been shown to occur in thermal treatments above 650°C and also for prolonged treatment at 610°C [39]. At temperatures above the melting point of Al, and under atmospheric pressure, SiC becomes thermodynamically

unstable and an invariant reaction occurs and leads to the formation of aluminum carbide (Al4C3) according to Equation 2 [40-41]

4Al + 3SiC → Al4C3 + 3Si (2)

The Gibbs free energy of formation of this reaction is large and highly negative between 900-

1200K which is in the range of processing temperature indicating that it is thermodynamically favorable.

The kinetics of Al4C3 formation has been investigated by Lloyd [42] who observed that its layer thickness, on SiC particles cannot be represented by a parabolic law for diffusion- controlled reaction. He proposed that aluminum carbide first nucleates at preferred sites on SiC,

SiC dissolution continues where SiC and Al are in direct contact, and finally dissolution of SiC occurs from areas separated by the interfacial reaction products. The dissolution of SiC depends upon the temperature, Si content in the melt, diffusion rate of Si in the melt and other such variables .Warrier and Lin[43] found that for short liquid state processing times, such as during infrared melting and infiltration, carbide growth was linear. These two results suggest that under suitable conditions immediately at contact of liquid aluminum and carbon, dissolution starts and carbide forms rapidly at the interface. The carbide layer thickens quickly, after which further carbon must diffuse through it to contact the molten aluminum and the reaction rate slows down.

The formation of Al4C3 in liquid or solid state processes is in the form of hexagonal platelets [44] that form discontinuously on the fiber surface and grow along a preferred direction.

The platelets will continue to nucleate and grow until they quickly form a continuous layer surrounding the fiber surface. Their continued growth is then determined by carbon diffusion in both solid and liquid state.

Even though this reaction enhances the wettability of SiC by the aluminum matrix [45], the formation of aluminum carbide is detrimental to composite properties. One of the biggest downsides to formation of this new interface is that it forms at the expense of the reinforcement.

For example, in the aluminum-Silicon carbide composite system, aluminum reacts readily to form (Al4C3), by diffusion of carbon out from the reinforcement to the particle and fiber surface, where it is dissolved by and reacts with aluminum [46]. This takes material away from the reinforcement, weakening it. In the case of silicon carbide or carbon fiber reinforcement, the reaction is uneven, so that the fiber develops narrow sections of concentrated stress [47].

Aluminum carbide grows as hexagonal plates, which will grow very large if the dissolution rate is slow relative to the diffusion rate of carbon through the interface, and these hard particles will impinge into the reinforcement. This will stress it and could cause cracks, to lower overall properties. The most damaging effect of reaction to composite properties may be in that the aluminum carbide layer is brittle. Considering toughness and ultimate strength of continuous fiber-reinforced composites, the brittle Al4C3 layer may be acceptable to some degree, as it permits energy absorbing fiber pullout and crack bridging. This layer is a source of crack initiation and propagation. It also prevents full transfer of load from the matrix to the reinforcement because of debonding of fibers and particles at lower loads, so that too much of it will lower overall mechanical properties in the fiber direction. Khan [48] demonstrated such effects by showing that as carbide thickness increased, a dramatic reduction in strength occurred.

As for the isotropic composite such as particle reinforced, it has been shown that any reaction phase helps to initiate and propagate cracks [49]. For these reasons, it is strongly desired to keep interfacial reaction as small as possible.

Because the carbide phase is detrimental to mechanical performance and forms so quickly, controlling and minimizing it are among the major concerns of liquid state processing.

Since wetting requires high temperatures, lowering the processing temperatures to reduce the reaction is not an approach. Processes which wet and cool quickly, such as rapid infrared process and squeeze casting, minimize the time of phase growth band are effective. Alloying additions are another option. Li et al. [49] demonstrated that titanium additions to the aluminum melt led to the formation of a thin layer of Titanium carbide (TiC) at the interface, which acted as a diffusion barrier for carbon to minimize aluminum carbide development. Lloyd et al. [41] showed that increasing silicon content in aluminum over seven percent can reduce the reaction for silicon carbide reinforcements. This occurs because the addition of silicon to the base alloy makes it to remain in equilibrium with silicon carbide at the processing temperatures.

In low silicon or no- silicon alloys, there has been another approach to control interfacial reaction, i.e. by coating. Coating has also been used to promote wetting of dispersions by a liquid alloy. Electroless copper and nickel coatings have been used widely in the past, but these coatings dissolve easily in the aluminum alloys and sometimes there is a formation of brittle intermetallic compounds impairing the ductility of the composite. Teng and Boyd[50] used the sol-gel process to coat Al2O3 or MgO on Silicon Carbide which restricts the interfacial reaction.

2.7 A356 casting alloy

This alloy was used as a matrix material for the composites in this study. This alloy is popular for casting operations because its silicon content permits it to be very fluid, and still yield a high strength value. The high hardness of silicon particles improves wear resistance. In the previous section it was reported that silicon minimizes interfacial reaction of silicon carbide to form aluminum carbide. The mechanism for this minimization is that since interfacial formation of aluminum carbide adds silicon in solution to the melt, the silicon in A356 acts to retard aluminum carbide formation. The properties of casting alloy A356 are outlined in table

2[51]. Phase diagram of Al-Si system [51] is shown in Figure 2.1. The Al-Si system has a eutectic reaction at 577ºC and a eutectic composition of 12.6 wt%. As aluminum and silicon solidify in different structures, respectively face centered cubic (FCC) and diamond cubic, two solid phases, α and β are produced. At high temperature, the hypoeutectic alloy forms a rich aluminum α- phase solid. The hypereutectic alloy forms almost pure β phase silicon. Very little silicon(1.65wt%) dissolves in the α phase aluminum and almost negligible amount of aluminum dissolves in the β phase. Application of A 356 alloy includes machine tool parts, aircraft wheels, pump parts, marine hardware, valve bodies etc [52].

Table 2 – Properties of A356 casting alloy[51]

Tensile Strength 221-262 MPa

Yield Strength 165-185 MPa

Elongation 5-6%

Elastic Modulus 72.4 GPa

Hardness 70-80 HB

3 Density 2.685Kg/m

º Liquidus Temperature 615 C

Solidus Temperature 577ºC

Linear Coefficient of Thermal 23.5µm/m•K at 20-300ºC

Expansion

Thermal Conductivity at 25ºC 151-155W/m•K

Figure 2.1: – Al-Si Phase Diagram [50]

2.8 Aging Behavior of Metal Matrix Composites

In metal matrix composites the difference in the thermal expansion coefficient ∆α, between the matrix and the reinforcement can be so large (as much as a factor of 4.7 between aluminum and SiC), that even a small change in temperature (during cooling from the fabrication temperatures or quenching) may generate residual stresses in the matrix[53]. The matrix can undergo plastic yielding as the magnitude of the residual stresses locally exceeds the yield strength giving rise to the formation of dislocations. The consequent development of dislocations in the matrix gives rise to a greater dislocation density in the matrix of the composite than in the unreinforced alloy. For example Arsenault and Fisher[54], employing transmission electron microscopy, observed a high density of dislocations (4 Χ 1012m-2) clustered around the ends of

SiC whiskers and forming low angle cell boundaries in a 20 vol% SiC whisker reinforced 6061 aluminum alloy. These dislocations can serve as heterogeneous sites for the nucleation of strengthening precipitates and can provide short circuit diffusion paths for solute atoms. It has been reported by I. dutta and Bourell [55] that the heterogeneous nucleation of the matrix precipitates on the dislocation lowers the elastic strain energy associated with the dislocations. In addition, it has also been suggested that the distortion field surrounding the edge dislocation reduces the energy barrier for nucleation of precipitates, efficiently increasing the aging kinetics.

As a result, both the nucleation and the growth of precipitates in the matrix can be drastically altered by the presence of reinforcements. Experimental results clearly show that, in a wide variety of precipitation-hardenable aluminum alloys, the alloy with brittle reinforcements exhibits a significantly shorter aging time to achieve peak strength than does the unreinforced matrix alloy at the same aging temperature. The faster development of strengthening precipitates

in the composite is commonly referred to as “accelerated aging” [53]. Figure 2.2 [56] shows the aging curve for the unreinforced 2124 aluminum alloy and 2124 aluminum alloy (powder metallurgy processed) with 13.2 vol % SiC whiskers, where the variation of Vickers microhardness is plotted as a function of aging time at a fixed aging temperature of 177°C.

Although the peak aging time for the unreinforced alloy is 10 to 12 hours, the peak hardness is reached in the composite after only 3-4 hours of aging because of the precipitation of metastable

S' Al2CuMg, which in the advanced stages of aging, is transformed to stable S precipitate. In addition in the study of aging behavior of 6061 Al alloy/SiC, Gupta and Surappa[57] revealed an increase in aging kinetics of the metallic matrix with an increase in volume fraction of SiC particulates. They attributed the accelerated aging kinetics to an increase in subgrain and grain boundary area in the metallic matrix. They noted a decrease in subgrain and grain size of the metallic matrix as a result of the presence of ceramic reinforcement. An increase in grain boundary area in the matrix will assist in increasing the frequency of nucleation of the strengthening phases as a result of the reduced activation barrier for the heterogeneous nucleation. Since precipitation is a diffusion aided process, both aging temperature and time have a strong influence on the precipitation kinetics of an age-hardenable alloy. In addition, in aluminum composites it has been found that peak aging time is very sensitive to temperature.

Chawla[58] examined the aging behavior of 2014 aluminum alloy and its composite (Al

2014+SiCp) as a function of time to reach peak strength vs. aging temperature. According to his observations, as the aging temperature increased, the peak aging time for the matrix of the composite decreased rather drastically. Furthermore, as the aging temperature decreased the difference in peak aging time between the matrix of the composite and the unreinforced 2014 alloy became smaller and at 150°C the difference in peak aging time is essentially negligible,

Figure.2.2: Variation of matrix microhardness as a function of aging time at 177°C for powder metallurgy processed 2124 aluminum alloy with and without 13.2 vol% SiC whiskers. (From

Christman and Suresh[55]

while it was 7 hours at 195°C.

2.9 Abrasive wear

Wear can be generally described as the removal of material from a surface in relative motion by mechanical and /or chemical processes [59]. There are four principal types of mechanical wear namely adhesive or galling wear, abrasive wear, corrosive wear and surface fatigue. In any particular instance of wear one may have any of these mechanisms operating either singly or in combination. In the latter case the situation may even be more complicated because of interaction between several mechanisms. For instance, the hard metal oxide formed on a steel surface by corrosion may then act as a fine abrasive to wear both of the surfaces. In this section we will restrict the discussion to abrasive wear.

Abrasive wear has been defined as the displacement of material caused by hard particles or hard protuberances where these hard particles or protuberances are forced against and moved along a solid surface [60]. Abraded Surfaces show damage which can range from fine scratches to deep gouges. A test that best visualizes this is where a pin of the material is worn against a disk or belt of abrasive, such as silicon carbide grinding belt. This particular test can isolate the wear to just abrasive as long as the pin is continually exposed to fresh abrasive. In this way, no other wear debris will build up to lubricate.

Abrasive wear can be classified as two body or three body abrasion. In the former, abrasive particles move freely over a material surface as in sand sliding down a chute or in the dredging of sand and gravel. In three body abrasion, abrasive particles act as interfacial elements between the solid body and counter body. This type of wear is important in particle reinforced composites as a composite pin is worn away, the particles get knocked out. If not removed from the wear track, they will remain to abrade against the pin, so that its wear volume increases

tremendously. Previous studies by Deshpande and Lin[61] have shown that this can make composite pins appear much less wear resistant than pure alloys if not accounted for. The discussion of the following parameters is worth, while reviewing abrasive wear [16].

2.9.1 Hardness: In general, increasing the hardness decreases the wear of a material, but there is no simple relationship between the two. According to the adhesive wear theory by Archard [60] wear volume is a function of sliding speed, normal load and material hardness. With the assumption that wear particles could be described as hemispherical particles of the same radius as the contact area, Archard[62] developed the following expression for wear rate, W(volume of material worn)

W=kdP/ 3H (3)

Where K= wear coefficient, d= sliding distance, P= applied normal load and H=bulk hardness of material. The theory predicted that enhanced wear resistance was associated with increase in hardness. However, this conclusion has not always been found to be valid. This simple relationship is obeyed only approximately by some pure metals, and most alloys show more complex behavior.

Abrasive wear studies reported by Hutchings[ 63]showed that if the hardness of the abrasive grit was more than 1.2 times that of the wearing surface then the abrasive medium would scratch the surface. Moore and King[64] also commented that the material‟s hardness determined the depth of indentation of the abrasive particles, thus influencing the relative penetration depth value. Wang and Rack[17] had reported that if this depth is greater than the critical indentation depth for fracture then the wear rate was high.

Larsen-Baswse and premartne[65] showed that craters are formed on a material‟s surface under conditions where the grit was trapped( but could still rotate) between the counterface and the specimen. Trapped grits that could not rotate resulted in microgrooving of material‟s surface.

These two different wear mechanisms were reported to be directly influenced by the ratio of hardnesses of the grit(Ha) and substrate material(Hm). In addition according to Burwell [66], in case of three body abrasion the hardness of the particle relative to that of each of the two rubbing surfaces will distinguish whether it will exist as loose rolling particles moving around between the two surfaces or whether it will become embedded in the softer of the two surfaces and act somewhat like a lap on the other surface. A common example is the situation of a journal bearing made of softer babbitt running in dirty surroundings where outside abrasive gets embedded into the babbitt bearing and serves to wear, not the softer babbitt , but the hardened steel shaft.

2.9.2 Effect of abrasive grit dimension: Sin et al. [65] indicated that the size of an abrasive grit had a direct influence on the associated wear mechanism. They characterized the effect of grit size by a ratio term, w/r, where w was the groove width and r was the radius of the spherical tip of the grit particle. Depending on this ratio, the abrasive particle would either plastically deform the surface or cut it. As the wear particle size decreased, the wear mechanism exhibited a transition from cutting to delamination wear.

For both types of wear (two-body and three body), a critical abrasive size has been observed [68]. The wear volume increases with abrasive grit size up to this critical dimension

[67]. At the critical grit dimension Larsen Badse[69] reported the maximum abrasive action in terms of the size of wear debris produced. Above this critical value the wear rate was largely independent of the abrasive grit size. The hardness difference between the abrasive particles and

wearing material influences the critical grit dimension [70]. Moore and Douthwaite [71] suggested that, during abrasive wear, the extent of the plastically strained region below an abraded surface depended on the abrasive grit size and the applied load. The overall depth of the plastic deformation was linearly related to the applied load and the grit dimensions.

2.9.3 Fracture Toughness: Above a critical loading the abrasive wear of brittle materials is controlled by the formation and/ or propagation of cracks. Values of fracture toughness of materials decrease to a crude approximation, with increasing hardness. The optimum combination of hardness and fracture toughness should be found in order to minimize wear in any particular tribosystem. Moore and King[62] cited that the rate of material removal and the associated wear process was determined by such factors as applied load, hardness of material and ratio of fracture toughness to material hardness. For conditions where this ratio was low, the wear rate was shown to be high, the debris are formed by a fracture mechanism. Hutchings[74] stated that aluminum based composites showed greater abrasive wear resistance, compared to unreinforced alloy, only under certain wear conditions. The crucial factor governing wear resistance was the extent of fracture of the reinforcing phase during the wear process. If brittle fracture was prevented then the material exhibited a low wear rate. For two -body abrasive wear

Zum Gahr [75]reported that abrasive particles initiated both microploughing and crack propagation of the wearing material only when the exerted load was above a critical value. He further proposed that fracture toughness of the material influenced the critical load and hence the wear resistance. Mathia and Lamy[76] studied fracture toughness on constituent phases by the application of a scratch test method, using increasing applied load. The phase was scratched with steadily increasing load. Using this technique the transition from ductile to brittle abrasion was

characterized by the onset of lateral cracking. Lateral cracking led to chipping. The ratio of fracture toughness to hardness was considered to be a determining factor in assessing the ductile or brittle nature of the material.

2.10 Wear Resistance of Metal Matrix Composites

For composite materials, the wear phenomena differ significantly from that of themonolithic metal due to the multiple phase nature of the composite. It is believed that during the initial stages of the abrasive wear test of MMCs, the abrasive particles from the counterface are capable of gouging the soft metal matrix from the polished composite surface, thus, protruding the reinforced particles to the composite surface. At this stage, the hard reinforced particles are in contact with the abrasive particles. The wear mechanism, from here on, is governed by various intrinsic factors such as hardness difference between counter faces, hardness of the reinforcement and the matrix, amount of reinforcement, the size and shape of reinforcement, hardness, distribution of reinforcement, bond strength between the reinforcement and the matrix, interfacial reaction between the reinforcement and the matrix, porosity and pore distribution.

A variety of mechanisms may be responsible for wear of composites. It is possible that when a composite wears individual phases may wear independent of each other. Indeed if this is true, the composite should obey, much like the density and the specific heat, the rules of mixtures. If the wear rate of the composite is controlled by the wear rate of the reinforcing phase, it is expected to substantially reduce the overall wear rate of the composite because the reinforcing phase is generally hard. On the contrary, if the reinforcing phase wears out faster than the matrix phase, which is unlikely, the overall wear rate of the composite should be about

the same as the matrix wear rate when the volume fraction of the reinforcing phase is low. When the volume fraction is large, (for instance more than 50%), it should be larger than the matrix wear rate. It is also conceivable that the reinforcing phase could be pulled out of the composite and trapped at the interface. In this event composite may wear by abrasive wear mechanisms. It is also possible that a completely different mechanism may operate in composites. For example, because the matrix reinforcing phase interface is generally weak, cracks may nucleate at the matrix/particle interface, either at the surface or below the surface. The surface cracks will be responsible for particle pull-out whereas the subsurface cracks may lead to delamination-type wear. In that case the wear rate of the overall composite depends, in a complicated way, on the particle size, interparticle spacing and the strength of the particle-matrix interface. As a result, the wear rate of a composite could be larger or smaller than the wear rates of the constituent phases.

For two-phase materials containing a hard, coarse second phase, Liou et al. [77] developed a relationship which determined the abrasive resistance. The treatment used by Liou et al [77] incorporated both micro-ploughing and brittle fracture mechanisms and reported the following expression for abrasive wear resistance, R(X):

R(X)αHeff = (1-X)Hm + αXHs (4)

Where Heff = effective hardness; X= Vol% second phase; Hm= hardness of the work hardened matrix phase: Hs= hardness of the second phase particles; α=critical ratio (Hm/Hs) such that α=1 indicates micro-ploughing and 0<α<1 indicates a brittle fracture mechanism. A lower critical ratio( which was strongly influenced by the work-hardening of the matrix) indicated a reduced abrasive wear resistance and a resultant lower effective hardness. Another expression to describe

the wear rate of a composite, Wc, based on the inverse rule of mixtures was given by [78]

fp (1 fal) {1 } { } { } (5) Wc Wp Wal

Where fp= area fraction of the particulate phase, fal= area fraction of aluminum phase Wp= particle wear rate, Wal=aluminum matrix wear rate. Generally this expression was considered too idealized given the assumption that each phase exhibited a specific and independent wear rate. A complete review of wear of metal matrix composites is not possible without laying emphasis on the following terms which play a crucial role in determining wear.

2.10.1 Nature of the reinforcement: Friction and wear behavior of metal matrix composites will depend on the nature of particle reinforcements and the matrix, and the interface between them. The particles can be softer or harder as compared to the matrix. The nonmetallic reinforcements dispersed in metal matrix composites can be divided into two different groups on the basis of hardness and wear mechanisms[79]: 1) hard particles with hardness 4-31Gpa, such as SiC, Al2O3, Silica, B4C and TiB2, and 2) soft particles with hardness below 4GPa, such as graphite and MoS2 which are primarily added for solid lubrication purposes. The hard ceramic particles appear to have a low adhesion to a metallic counterface. The asperity of the counterface can easily plow through softer particles like graphite while it cannot do so for harder particles like alumina or silicon carbide. Figure 2.3 shows the variation of specific wear rate as a function of volume fraction in various systems [79]. Char is a soft porous particle containing carbon and other hard mineral oxides, and the composites containing char show higher specific wear rate.

Alumina is harder than char, but softer than silicon carbide, and aluminum-alumina shows a higher specific rate up to 5 vol% than Al-SiC composite. However, Al-Al2O3 composites show a

far lower specific rate than Aluminum-Char composite. Roy et al. [80] also have demonstrated that composites containing SiC, TiC, TiB2 or B4C particles exhibit lower wear rate when compared to base material. However, aluminum alloy base composites containing graphite has the lowest wear because graphite is not only soft but shears easily along the basal plane of its hexagonal close packed lattice in suitable environment and acts as a solid lubricant. The composites containing solid lubricant-like graphite have low wear due to its transfer on the tribosurfaces and formation of a film of graphite between the composite and the counterface.

However, higher graphite contents increase the extent of cover of the graphite film on the alloy sliding surface and there is little change in the wear rate. Once graphite build-up on sliding surfaces becomes thick enough, delamination wear ensues, resulting in increased wear [81].

Figure.2.3 Specific wear rate in several aluminum base particulate composites sliding against steel as a function of particle volume fraction[79]

Thus it may be concluded that the harder the reinforcement particles the lower will be sliding wear in a composite, with the exception of solid lubricating dispersoids of the right crystal structure which impart low wear rate despite their low hardness.

2.10.2 Effect of Reinforcement : It has been observed in many investigations that reinforcing the aluminum alloys with SiC or Al2O3 improve the wear and abrasion resistance. According to the reports submitted by Zum Ghar [82] and Hutchings [83], aluminum matrix composites usually exhibit isotropic abrasive wear behavior and the wear resistance of these composites is greater than the virgin alloy as long as the reinforcing particles remain intact on the worn surface. During Abrasive wear, the overall wear depends on the ability of the abrasive to penetrate into the specimen surface and the extent of penetration. In the case of alloy the depth of penetration is governed by the hardness of the specimen surface, abrasive size and applied load.

But in case of composite, the depth of penetration of the abrasives is primarily governed by the protruded hard ceramic reinforcement. Thus, the major portion of the load is carried by the reinforced particles and the penetration of abrasive is restricted by these hard reinforcements. If the effective load on the individual particle increased above its fracture strength, the particles get fractured. The depth of penetration of the abrasive may also depend on the dimension of the matrix surface (inter-particle distance) in composite, which is free from reinforcement [84]. This is because of the fact that the abrasive will not get any hindrance from the reinforcing phase to penetrate into these regions of the composite. In view of this fact, it is expected that uniform distribution of the particle and finer-inter particle distance will lead to high wear resistance.

The improvement in abrasive wear resistance in aluminum matrix composites could be

described in terms of relative wear resistance (RWR) which is defined as the ratio of wear resistance of composites to that of the alloy. The RWR of aluminum matrix composites has been reported to be varied between 1 and 10 depending on the wear condition. In case of low abrasive wear (Rubber wheel abrasion test), the RWR is high enough. In these tests the stress developed during the abrasive-particle interaction is not enough to exceed a limiting stress required for local elastic deformation of rubber wheel to engulf the abrading particle. However in case of high stress abrasive wear characterized by pin-on-disc arrangement, the RWR was reported to be less

[85-86]. It is observed that under higher applied load and coarser abrasive size the value of RWR may be less than unity. This is primarily attributed to fracture and fragmentation of particles and their removal from the wider and deeper wear grooves generated on the specimen surface under such severe wear condition.

With reference to dry sliding wear of A356 reinforced with SiCp Pramila Bai [87] observed that with increasing applied pressure the wear behavior of the unreinforced alloy was dominated by extensive plastic flow of the alloy surface and significant wear debris formation.

The addition of SiC reduced the wear for the applied pressure range examined. Silicon carbide particles were reported to minimize this plastic deformation on the wearing surface and promoted the formation of an iron-rich layer on the composite surface acting as a lubricant and thereby reducing wear. It was observed by Das [88] that at the onset of severe wear with increasing the applied load and sliding velocity the bulk temperature in the aluminum-7% silicon alloys exceeds 200°C. Reinforcing the matrix with SiC particles minimized the friction induced surface heating and seizure resistance. The particulate (or whisker) reinforcement of aluminum alloys is known to improve the thermal stability of these materials [89]. The strength of the composites is maintained to a higher temperature than that of the base alloy. So, it can be

suggested that the retention of room temperature strength (and ductility) at elevated temperatures may be the reason for the suppression of severe wear in the SiC reinforced composites.

2.10.3 Effect of Increasing volume fraction of reinforcement on wear: The abrasion wear resistance of the composites increases with the reinforcement volume fraction if the damage in the reinforcements, cracking or pull-out of the matrix, is negligible [90]. At first glance, the experimental results obtained up to now on the abrasive wear behavior of metallic matrix reinforced materials are contradictory. Some authors[91-92] have reported a monotonic increase in the wear resistance of the material with increasing volume fraction of reinforcement particles, while other authors[93] observed the opposite: the wear resistance decreases continuously with increasing volume fraction of reinforcement particles. In the study of Al6061 alloy reinforced with Al2O3 fibers Wang and Hutchings[90] reported that for small abrasive particles, the wear resistance of the composite increased with increasing fiber volume fraction. However, as the abrasive particle size increased, a transition in wear behavior was observed. With the larger abrasive particles, the wear resistance decreased with increasing fiber volume fraction (≥20 vol%) and the worn surface revealed fiber fracture and extensive debonding at the fiber-matrix interface. This increase in fiber volume fraction resulted in a reduction in the extent of plastic deformation in the matrix and increased the probability of fiber abrasive interaction. This led to a greater material removal by fiber fracture and then micro-cracking. Therefore, the extent of fracture of the reinforcing phase played a critical role in determining the wear behavior of the composite.

Hutchings[82] and Zum-Gahr [83] have reviewed the tribological behavior of this type of materials and proposed a number of mechanisms for its mass loss under abrasive wear

conditions. According to the analysis carried out by these authors, the wear behavior of metallic matrix reinforced materials depends, first, on the relation between the size of the reinforcement particles and the dimensions of the damage caused by the abrasive (for example, the width and depth of the grooves). If the damage is much larger than the size of the reinforcement particles, the metal matrix reinforced composite behaved as a homogenous material. In this situation an increase in the proportion of reinforcement particles will lead to an increase in the wear resistance of the material, due to the corresponding increase in its hardness, provided that the wear mechanism does not change. Conversely, if the reinforcement material is more brittle than the matrix, and if the abrasive is harder than the reinforcement material, increasing the proportion of reinforcement particles can lead to a change in the wear mechanism from microploughing to microcutting, resulting in a decrease of the wear resistance of the material. If the dimension of the damage caused by the abrasive is similar to or smaller than the size of the reinforcement particles, the matrix and the reinforcement will respond individually to the action of the abrasive, according to their different mechanical properties. In this case the relation between the hardness of the reinforcement particles and the hardness of the abrasive will determine the wear behavior of the material. If the hardness of the abrasive is lower than that of the reinforcement particles, the matrix will be damaged preferentially. If the hardness of the reinforcement particles is lower than that of the abrasive and the loading conditions are severe, the particles will tend to fracture and a sudden increase in wear would occur. In such a situation increasing the amount of reinforcement would lead to an increase in wear rate. Studies undertaken at low and mild loads indicate that increasing the volume fraction of the reinforcement improves wear resistance.

2.10.4 Interfacial Strength: Wear behavior will depend on the strength of the interface between the reinforcement particles and the matrix since, if the interfaces are weak, the reinforcement particles can easily be pulled out by the action of the abrasive, leading to considerable material loss. The matrix–reinforcement interface has a strong influence on thermo-mechanical properties of hard phase containing material. The properties of the interface depend on physical and chemical properties of the matrix and hard particles, and on the reaction between these materials during consolidation and subsequent heat treatment. It has been reported that reaction layers, which are relatively thick (around 1µm, weaken mechanical properties of the composites, while the layers thinner than 0.1 µm might be beneficial [94]. Wang [95] corroborated the above fact in his observation that that Ti50Ni25Cu25 particle reinforced aluminum matrix composite with interfacial reactions exhibits lower wear resistance than the composite without interfacial reactions. In the rubber wheel abrasion test (RWAT), Bhansali and Mehrabian[96] reported that Al alloy reinforced with SiCp showed similar wear resistance to MMC‟s reinforced with same proportion of alumina particles of the same size. SiC being appreciably harder than alumina was expected to give higher wear resistance. Microstructures showed that weakness of the interface with the matrix played a major role in preventing the reinforcement from achieving its full potential.

Prasad and calvert[97] demonstrated the influence of interfacial bond in the three body abrasive wear study of polymethyl methacrylate(PMMA) reinforced by quartz and glass beads.

SiC, SiO2 and CaCO3 abrasive slurries in water were used in wear tests. A silane treatment of the fillers was carried out and the results indicated that silane helped in reducing composite wear by improving the particle-matrix bond so that the matrix supports the particle at the boundary and reduces the tendency of these regions to chip off when loaded. With reference to dry sliding wear, Modi [98] showed that reinforcement/matrix interfacial strength and the dispersoid shape

both greatly influenced composite wear rate. Composites reinforced with SiC particles revealed a lower wear rate than those containing SiC fibers. The SiC fibers were associated with poor interfacial bonding that promoted dispersoid pullout, leading to void formation, a higher probability of crack nucleation at these voids and higher wear rate due to three-body abrasion effects.

2.10.5 Shape of the dispersoid and orientation: The shape and the aspect ratio of the dispersed phase are very important in determining the load bearing capacity. Accordingly, the composites have been classified as a) Fiber-Reinforced composites with both continuous and discontinuous fibers and b) Particle or whisker reinforced composites. In discontinuous fiber composites or particle reinforced composites the load is transmitted to the dispersoid through the matrix.

Mismatch of strain in the matrix and the dispersoid across the interface results in a shear stress as described in simplistic “shear lag models for discontinuous fibers and particles in a matrix. If the shear stress that develops at the interface exceeds the strength of the interface there will be debonding at the interface and this can develop into a crack that can propagate. For a given condition of loading, the load shared by the dispersed phase increases with its aspect ratio.

Ma et al [99] examined the abrasive wear behavior of aluminum containing different reinforcement phases. High stress abrasion tests(pin-on-disc) were performed on an Al 6061 alloy reinforced with SiCw, SiCp and SiCf using emery papers (20-28 and 14-20µm). The abrasive wear resistance was found to be dependent on the reinforcement type and on the properties of the matrix alloy. He reported the abrasive wear resistance of the composite in the following order:

(low wear rate) SiCf/6061>SiCw/6061>SiCw/6061+SiCp/6061>SiCp/6061(high wear rate)

Orientation of the fibers also has an influence on the abrasive wear behavior of the composites. The orientation of fibers may vary between parallel and perpendicular to the wearing surface. Fibers parallel to the surface may be dug out more easily than those perpendicular to it.

This has been verified by [100-101] who reported that the abrasive wear loss increased by changing the fiber orientation from perpendicular to parallel to the wearing surface.

2.10.6 Particle size: According to Hutchings[74] the dimensions of the contact area of each abrasive or erosive particle, compared with the size of the reinforcement, will determine whether the composite responds heterogeneously to the particle contact or whether it acts effectively as a homogenous solid. In the first case as shown in the fig.2.4 (a) the particle contact zone is appreciably smaller than the reinforcement size, and the individual elements of the reinforcement can be considered to act independently. Under these conditions a relatively high wear resistance compared with that of matrix may result if fracture of the reinforcement is avoided. However, if the reinforcement size is much smaller than the scale of the contact zone as shown in fig. 2.4(b), even if the reinforcement phase does not fracture it will simply be removed by the plastic processes (e.g., ploughing and cutting) along with the matrix, and no major benefits of the reinforcement can be expected, apart from those associated with the modest increase in the bulk hardness due to reinforcement. Alpas[102] observed that when the particulate size was of the order of the surface roughness of the steel counterface(e.g. 1-2µm), the reinforcement could provide only limited protection to the matrix. Damage in the form of particulate fracture, pull- out, and the plastic deformation of the matrix were observed at low loads for the alloy containing

2.4µm SiC. In contrast, in the alloys containing larger (15.8µm) SiC particulates, the reinforcement was able to carry the applied load and protect the Al matrix .He proposed eq. (6)

Figure2.4: Drawing illustrating the concept of dimensions of contact area. (a) The composite simulates a heterogenous material (b) With a finer microstructure, the wearing material acts as a homogeneous material [74]

in respect to the wear of particulate-reinforced MMCs based on the delamination theory.

V cLS cLSVf (6) (1 Vf )d

Where λ = (1- Vf )d/Vf, c is a constant dependent upon material properties, L the normal load, S the sliding distance, Vf the particle volume fraction and d the average particle size. Eq.

(6) states that the wear volume is not only proportional to the sliding distance and applied load, but also increases with the particle volume fraction and decrease with particle size. The particle size of the reinforcement determines the load bearing capacity of the reinforcement and the surface area for bonding i.e. interfaces between the constituents of the composite. In their study of SiC particle reinforced aluminum alloy composites, Sahinet [103] observed that increasing the particle size of SiC resulted in increasing density of the composite, thus, resulting in decreasing wear rate. The decrease in the wear rate with increase in the particle size was also evident in the study of Al2O3 – reinforced aluminum composites by Yilmaz et al. [104].This theory was further supported by Chung and Hwang[105] who stated that for a constant SiCp volume fraction,

MMC‟s containing coarser particles exhibited higher wear resistance. Delamination wear was considered the predominant wear mechanism. The coarser SiCp were reported to provide greater resistance to the propagation of subsurface cracks as compared to finer SiC fractions. The findings presented by Chung and Hwang are particularly significant since the use of aluminum as the matrix avoided the effects of ageing and formation of intermetallic compounds within the microstructure. Kok[106] attributed the higher wear resistance of the composites containing

66µm Al2O3 particles compared with those containing 16 µm Al2O3 particles to the change in wear mechanism from cutting and ploughing as the particle size increases under sliding

conditions. He observed that these large particles very effectively resist the penetration and cutting into the surface, and are not easily cut out by the slider because of their large size high values of hardness and good bonding with the matrix.

3.0 Experimental

3.1 Materials:

The as-cast ingot A356/SiCp used in this study was purchased from Duralcan.

The weight percentage of the alloying elements of A356 according to AMS standard is given in Table 2.Silicon is the main alloying element; it imparts high fluidity and low shrinkage, which result in good castability and weldability. The high hardness of the silicon particles also improves wear resistance. For the reference purpose an alloy of aluminum with 10% silicon by weight was also taken.

Table 3. Nominal Composition of A356 aluminum alloy (wt %)

Si Cu Mg Mn Fe Zn

6.5 to 0.25 0.2 to 0.35 0.6 0.35

7.5 max. 0.45 max. max. max

3.2 Chemical Analysis of Composite:

To determine the content of SiC in composite, a thin section of the composite was cut using a slow speed diamond saw and was polished on both sides to ensure that the oxide layer is removed and was weighed on the analytical balance. It was then dissolved in diluted HCl and the SiC particulate was filtered out on a filter paper after the complete reaction of aluminum with HCl as shown below:

2Al + 6HCl 2AlCl3 + 3H2 (7)

The filter paper was kept in a furnace at 80°C for 12 hours until it dried out completely.

The filtered SiC particulate was carefully weighed and the volume fraction of SiC was calculated using the following formula: 42

WSiC

SiC VSiC (8) WSiC WAl

Sic Al

where WSiC is the weight of SiC particulate, Wal is the weight of aluminum, ρSiC is the density of

SiC particulate and ρal is the density of aluminum. Weight of aluminum was calculated by subtracting the weight of SiC particulate from the initial weight of the composite. This method gives approximate values of the particulate content since the volume of the alloying elements and pores is not considered.

3.3 Particle size measurement:

For calculating particle size six micrographs of the composite showing a uniform distribution of particles at 400X using optical microscope were taken. But this was not the actual magnification as it did not include the magnification of the camera lens. In order to calculate the actual magnification a scale micrometer having the line spacing of

10µm was used and actual magnification was found out by measuring the spacing between the lines after obtaining the micrographs of scale micrometer. For measuring the particle size around 30-40 particles were considered and their dimensions were measured and converted to the actual length by using the magnification result. This calculation gave a range existing from smallest to largest particle size.

3.4 Density Measurement:

The densities of the as cast Al/SiC composite and Al-Si(10wt%) alloy were measured using Archimedes principle of water displacement. In this experiment a copper wire was taken and its weight in air (W1) and water (W2) was measured. The samples of Al-Si alloy and Al/SiC were attached to the wire and their weight in air (W3) and water (W4) was measured. After that, following measurements were done in order to calculate the final density of the alloy and the composite.

(a) Sample weight = (W3-W1) = Wa

(b) Wire buoyancy = (W2-W1)= Wb

(d) Sample buoyancy only = (Wc-Wb)= Wd

(e) Sample Volume (V) = Wd (9) Dw

Where Dw is the density of water

Finally sample density (d) was calculated dividing sample weight (Wa) by sample

Volume (V)

Wa Sample density (d) = V (10)

Four specimens of alloy and composite were taken and final density was calculated by taking the average of these.

3.5 Wear Test:

The wear test was carried out with a pin-on-disk tester as shown in Figure 3.1.

The tests were carried out at room temperature. Each specimen was ground up to grade

600 abrasive paper to ensure that the wear surface was in complete contact with the

2 abrasive counterface. Rectangular specimens having contact area 16.129 mm were loaded against a disk, which rotated at 184 rpm. A 60 grit size (268 microns particle size) was bonded on the circular disc with the help of a double- sided scotch tape so that the emery paper was rigidly fixed on the disc surface and the assembly acted as an abrasive medium. The specimens were kept stationary, while the disc was in rotatory motion.

Thus, these particles acted like micro-indenter and machining tools as well, under applied load and rotatory motion. The applied normal loads used were 1.625N, 2.166N, 2.714N,

4.17N, 6.2 and 10.38N. Thus the applied pressure varied from 0.1049MPa to 0.67MPa.

Each test was carried out on a new emery paper. The sliding distance and velocity were kept constant at 80 meters and 1.54m/sec respectively for each specimen. The samples were cleaned prior to and after each interval of wear test with acetone. The wear rates of the composite and alloy were calculated by measuring the difference in weight of the specimens measured before and after the tests (measured with an analytical balance

Mettler AJ100, Hightstown, NJ of 0.1mg precision) per unit density and sliding distance(80m). The unit of the wear rate came out to be mm3/m. For each test condition, three tests were performed and the average was used.

Figure 3.1: Schematic representation of a pin-on-disk wear tester

3.6 Heat Treatment:

The heat treatments used in this study to increase the hardness and wear resistance of the composite were three step processes:

Solution treatment of the composite which resulted in the dissolution of the

soluble phases was carried out at 565°C for 2 hours.

Quenching which resulted in the development of a supersaturated state was done

by rapidly cooling the specimen in cold water.

Artificial aging which involved the precipitation of solute atoms was carried out

immediately after quenching at 180°C.

3.7 Hardness Testing:

The hardness of the composite and its aging response was monitored using

Rockwell B hardness tests using 1/16" diameter steel ball with a 100-kilogram load.

Before each hardness test series, the specimens, which were of thickness 8mm were metallographically polished with a 240 grit SiC paper until the oxide layer was removed and the opposite sides were perfectly parallel. Five specimens were used and an average of five measurements was reported.

3.8 X-Ray Diffraction:

In order to find out new precipitates or phases, X-ray diffraction was carried out using a pan analytical Philips system. This system had a CuKα anode (λ=.154065nm) and

47 a 2kW sealed tube operated at 45kV and 20mA. The step size and step time was 0.01 and

1second respectively. 1/8th inch size slit was used for the source.

3.9 Microscopic Examination

Composite (Al/SiC), alloy (Al-Si) and worn samples were prepared for evaluation by optical and scanning electron microscopy. Cross sections of the as cast composite, alloy and of the samples at each heat treatment were cold mounted in epoxy, since the heat of hot mounting the samples could affect the microstructure. These samples were rough polished using silicon carbide emery paper and then fine polished using Alumina paste down to 0.05μm finish. Mounting and polishing of the worn samples was not done in order to identify actual wear mechanism. Samples were not etched so that the volume percent of the reinforcement was properly represented.

The mounted cross sections of composite, alloy and worn samples were sputter coated with gold and palladium in order to minimize charging because of the presence of

SiC which is a non conductor for scanning electron microscopic examination. The microstructures were examined using a Philips XL30 Environmental Scanning Electron

Microscope with a field emission gun electron source. For finding out the composition of the precipitate energy dispersive X-ray analysis was also performed.

4.0 Results

4.1 Composite:

The microstructure of the as received composite shown in Fig. 4.1a shows a reasonably uniform distribution of the SiC particles dispersed in the aluminum matrix.

The matrix seems to be well bonded with no visible voids. A magnified view of the microstructure showed in Fig. 4.1b shows a SiC particle embedded in aluminum matrix.

(a)

(b)

Figure 4.1: Microstructure of the composite showing (a) uniform distribution of SiC particles in the matrix and (b) SiC particle embedded in aluminum matrix

The Volume fraction of SiC calculated by the chemical analysis was found to be between

23% which agreed with the composition provided by the vendor. The SiC particle size which

was calculated with the help of scale micrometer was found to be between 5µm-20 µm. Note

that the sample was not etched so that the volume percent of the reinforcement is properly

represented. The densities of the as cast Al/SiC composite and Al-Si(10wt%) alloy were

measured using Archimedes principle of water displacement were found to be 2.7543gm/cc

and 2.6136gm/cc respectively.

Fig.4.2 (a) Optical Micrograph of the composite at 1500X (b) Optical Micrograph of the scale micrometer at 1500X

4.2 Aging Studies

Figure 4.3 shows the variation of composite hardness (Rockwell B) as a function of aging time at a temperature of 180°C. The composite reached peak hardness after aging for 4-6 hours. The hardness increases initially with increasing time. It reached a maximum at around 5 hours. Further aging reduced the composite hardness.

SEM micrographs of the as cast composite and the heat treated showing precipitates aged at 180 °C for various times are shown in fig. 4.4 . The micrographs are in agreement with the aging curve. The peak aged condition as shown in fig 4.4b is characterized by an optimum precipitate size and a large number of precipitates as is evident from the micrograph. As the aging time increases the precipitates become coarse and the distance between them becomes large. This condition corresponds to fig. 4.3c and fig. 4.4d in which the composite was aged for 12 and 40 hours respectively. No precipitates were observed in the as- cast composite. For reference purpose, heat treatment of the as-cast

Al-Si(10wt%) alloy was also carried out. Figure 4.5 shows the micrographs as cast and heat treated alloys. The heat treatment of the alloy was carried out for the same time and temperature as the Al/SiC composite.

Table 4. Rockwell B hardness for composites and alloys used for wear tests(Average

value of 5 hardness readings)

Material Aging time(hrs) Hardness(Rockwell B)

As Cast Al/SiCp ___ 57.53

Al/SiCp 5 73.9

Al/SiCp 12 67.26

Al/SiCp 40 58.76

Al-Si(10wt%) ___ 14.5

Al-Si(10wt%) 5 52.1

Al-Si(10wt%) 12 44.8

Al-Si (10wt%) 40 28.6

Rockwell Hardness(B) VS Aging time (hrs)

Rockwell Hardness(B) Rockwell 20

0 0 10 20 30 40 50 Time (hrs)

Fig. 4.3. Variation of Rockwell hardness with Aging time

(a)

(b)

Fig. 4.4. Micrographs of the Composite showing (a) As cast Composite (b) Composite aged for 5 hours

(c)

(d)

Fig. 4.4.(Continue) Micrographs of the Composite showing (c) Composite aged for 12 hours and (d) Composite aged for 40 hours

(a)

(b)

Fig. 4.5. Micrographs of the Alloy showing (a) As cast Alloy Al-Si(10wt%)

(b) Alloy aged for 5 hours

(c)

(d)

Fig. 4.5.(Continue) Micrographs of the Alloy showing (c) Alloy aged for 12 hours and (d) Alloy aged for 40 hours

4.3 X-ray Diffraction

Fig. 4.6: X-ray diffraction pattern of the as-cast and aged composite

4.4 Energy dispersive X-ray analysis of precipitate

Fig. 4.7 (a) Micrograph showing a precipitate (b) EDS analysis of the precipitate

Table -5 EDS result of precipitate present in the composite after aging

Element Weight % Atomic %

Carbon 50.93 69.5

oxygen 3.36 3.45

Aluminum 16.09 9.78

Silicon 29.62 17.28

The composition of the precipitate according to the EDS analysis is shown in Table 5 in the results section. According to it the precipitate is composed of 69.5 at% carbon,

3.45 at% oxygen, 9.78 at% aluminum and 17.28at% silicon. Oxygen might have been introduced during processing or polishing. But according to the phase diagram shown in figure 2.1 no such intermetallic phase of aluminum and silicon exists. Neither does X-ray diffraction shows the presence of any new phase. It is possible that the beam might have penetrated into the precipitate and picked up aluminum and carbon from the matrix and

SiC particles respectively.

4.5 Wear Behavior

Figure 4.7 represents the variation of wear rate of the composite and the alloy for three successive tests as a function of applied load for a fixed sliding distance of 80m against a 60 grit SiC abrasive paper. Figure 4.8 represents wear rate by taking the average of three wear tests. It may be noted that wear rate of the alloy and the composite increased with increasing applied load. The Al/SiC composite has a minimum value of wear rate at peak-aging (5hrs aging) condition corresponding to the maximum value in 60 hardness. Aging for 12 hours resulted in lowering the wear resistance of the composite.

Further aging (40hours) resulted in more wear rate. However, as compared to the as-cast composite aging treatment resulted in an improvement in wear resistance. It was also observed that at the highest pressure of 0.67 MPa, the wear loss in the composites aged for 12 and 40 hours was very close to the as cast composite. Also on all loads it was observed that the wear rate in the composites was much less than the as cast Al-Si

(10wt%) alloy. For reference purpose, wear rate of as-cast Al-Si(10wt%) alloy and heat treated alloys was also compared. Time and temperature for aging was kept similar to that of Al/SiC composite. Heat treatment resulted in an improvement in wear behavior of the alloy as compared to as cast alloy. However the variation of wear behavior with aging time which was observed in case of Al/SiC composite was not observed. In the case of alloys, wear rate was almost similar for the alloys aged for 5, 12 and 40 hours and no variation among them was observed. The wear rate of the aged alloys was found more than the as cast and aged composites.

4.6 Worn Surfaces

To understand the wear mechanism of aluminum alloy and its composite, SEM observation of the wear surfaces was done. Scanning electron micrographs of the wear surfaces showing microcutting and microcracking are shown in figure 4.12 and figure

4.13 respectively. Microcracking was seen in all specimens. However this phenomenon was more predominant in composites as compared to the as cast Al-Si (10wt%) alloy.

Figure 4.11 shows the worn surface of the Al/SiC composite after the wear test. As compared to figure 4.1a not much SiC particles were seen on the surface of the composite after the wear test. 61

Wear data for Al/SiC Composite and as cast Al-Si alloy

Fig. 4.8. Variation of Wear Rate with Pressure at a sliding distance of 80m

Wear graph for Al/SiC Composite and As Cast Al-Si alloy

Fig. 4.9. Variation of Wear Rate with Pressure taking the average of three tests at

Sliding distance of 80m

Wear Data for as cast Al-Si and aged Al-Si alloys

Fig. 4.10 Variation of Wear Rate with Pressure at a sliding distance of 80m

Figure 4.11: SEM Micrograph of the 12 hours aged composite surface after wear

(a)

(b)

(c)

Fig 4.12 SEM Micrographs showing microcutting in (a) Alloy (b) As cast Composite

(d)

(e)

Fig 4.12. (Continue)SEM Micrographs showing microcutting in (d) Composite aged for

12 hours (e) Composite Aged for 40 hours

(a)

(b)

(c)

Fig. 4.13. Micrographs showing microcracking in (a) alloy (b) As cast Composite (c)

Composite aged for 5 hours

(d)

(e)

Fig. 4.13. (Continue) Micrographs showing microcracking in (d) Composite aged for 12 hours and (e) Composite aged for 40 hours

5.0 Discussion

5.1 Aging Studies:

As it can be seen from the phase diagram of Al-Si system (figure 2.1), the equilibrium solid solubility of silicon in aluminum increases with increase in temperature.

At around 565°C the solubility of silicon is around 1.6 wt%. When the composite is kept at this temperature for two hours to allow sufficient diffusion, part of the silicon will be taken completely into the solid solution. When the composite is quenched in water from such a high temperature, then the solid solution becomes supersaturated with silicon and as a tendency to achieve thermodynamic equilibrium precipitation of silicon takes place on subsequent aging.

This increase in hardness due to aging can be explained by dislocation theory. A precipitate particle acts as an obstruction to the motion of a dislocation. Such an obstruction provides resistance to the motion of dislocation and hence increases strength and hardness. For a dislocation to move it must either cut through the precipitate particles or move between them. In both the cases an increase in stress is required as compared to the matrix which does not contain precipitates. This additional stress calculated by

Orowan is given by equation (11). He assumed dislocations to bend in the form of expanding loops around the precipitate particles.

τ = 2Gb/d (11)

Where τ is the shear stress, d is the distance between two particles, and b is the

Burgers vector. Thus smaller the distance, d, between precipitates, the harder the material is. This explains the reason for the increase in hardness with the increase in the number of

70 precipitate particles corresponding to the peak aged condition. Beyond this stage as can be seen from the figure 4.4 the precipitates begin to grow and become less in number.

The smaller precipitates will dissolve and the larger ones will grow. This coarsening of microstructure is called “Ostwald Ripening “The driving force is the decrease of surface area and therefore, surface energy. For the smaller ones to dissolve and the larger ones to grow, there must be diffusion of the solute. This can be more readily understood from the from the free energy vs. composition curves for the α and β phases in the Al-Si system.

Small β (Si) precipitates have a higher energy than the larger ones because of more surface area for the same volume. For small β precipitates the tangent line connecting the curves for α and β is tangent to the α curve at higher B concentrations as compared to the larger β precipitates. This in turn produces a concentration gradient in α aluminum matrix between small and large β precipitates. This concentration gradient causes the solute to diffuse in the direction of large precipitates away from the small, so that small precipitates shrink and disappear while larger precipitates grow. As the precipitates become large and few, distance d between the precipitates increases, making it easier for the dislocations to pass through and resulting in a decrease in shear stress and therefore hardness.

β- small diameter particle

β-large diameter particle

Free Energy Free

Cα1 Cα2 Composition, % B 1 Cα1

Figure 5.1 Free Energy vs. Composition diagram for α and β phases in Al-Si system

5.2 X-Ray Diffraction analysis

X-Ray diffraction was done in order to detect the presence of any new phase in the composite after aging. As can be seen from the X-Ray scan (figure 4.6) that peaks of aluminum, silicon and silicon carbide were detected in both the as cast and the aged composite. No new peak was observed in the aged composite. This may be due to the fact that the new phase formed (precipitates) was of silicon, which was also present in the cast composite in form of small precipitates as can be seen from the worn surface in figure

4.10b. These precipitates were formed due to natural aging and were too small to be detected at 1250X in figure 4.4(a). This may be the reason for the increase in silicon peak intensity in the aged composite. There was possibility of forming Mg2Si precipitates, but the concentration of Mg was very less in the alloy (0.2-0.4 wt %) and with this concentration it could not get supersaturated and will be stable in the α phase itself.

However there was possibility of formation of aluminum carbide at the interface of aluminum matrix and silicon carbide particle. The thermodynamic analysis of its formation is done below.

SiC(s) Si(s) + C(s) (12)

ΔG= 58520 + 5.434 logT – 23.74T Joules/mole[107]

Al4C3(s) 4Al(s) + 3C(s) (13)

ΔG = 215688 – 4.18T Joules/mole [107]

Multiplying equation (10) by 3 and subtracting equation (11) will give

3SiC(s) + 4Al(s) Al4C3(s) + 3Si (s) [107] (14)

At 838K, which is the solutionizing temperature of the composite used in the study, free energy change for the above reaction (14) will be

ΔG = - 26.98KJ/Mole which is negative and the reaction is expected to go in the forward direction according to thermodynamics. But this value of ΔG was calculated assuming that the activity of silicon was one, which is however unlikely since the matrix is already supersaturated with silicon. The maximum solubility of silicon in aluminum is 1.65 wt% and the matrix contains around 6-7 wt% of silicon. The excess silicon in the matrix makes it remain in equilibrium with SiC and the tendency of the reaction to go in the forward direction is reduced according to the Le Chatelier’s principle. However formation of small amount of aluminum carbide cannot be ruled out.

5.3 Wear Studies

5.3.1 Effect of load on wear

As it was seen in the last section (figure 4.8, 4.9 and 4.10) that wear rate of each specimen increased on increasing the load following almost a linear relationship. The abrasive particles will act as microindentors and an increase in load will lead to an increase in the depth of penetration of these abrasives into the sample. This would in turn result in gouging off the matrix from deeper areas resulting in more wear loss. Moreover, an increase in load would also increase the depth of scratches made on the specimen surface. Wear studies by Wang and Rack [108] also suggest that wear behavior of the composite would depend upon the ratio of depth of groove (i.e. penetration depth made by the abrasive) to the particle size of the composites. They further proposed that above the value of unity for these ratios, the particulate reinforcement would give only limited

74 protection to the matrix. In the present study, when compared to figure 4.1b not much

SiC particles were seen on the surface of the composite after the wear test as is evident from figure 4.11. This finding suggests that the depth of the scratches was greater than the vertical dimension of the particles in most cases. When the applied stress exceeds a critical value, matrix cracking would occur as can be seen from figure 4.13. Matrix yielding was observed in as cast Al-Si(10wt%) alloy but not in Al/SiC composite.

Further, nucleation and propagation of these cracks would lead to the formation of coarse debris enhancing wear loss. Fracture of the SiC particles would occur when the applied load produces stresses greater than the fracture strength (3440MPa) of the SiC phase.

Fractured particles also tend to come out of the wear surface. Because of the above mentioned reasons, the alloy as well as the composite suffers from higher wear rate at higher applied loads. Before discussing the wear behavior observed in the composites a brief discussion on the mechanisms of matrix and particle removal based on the observations in the present study is a pre-requisite.

5.3.2 Mechanism of Matrix Removal

As is evident from the figures 4.12 and 4.13 in the previous section , the worn surfaces of the as cast and the aged composites showed microcutting and microcracking wear. Microcutting results in the formation of grooves. The depth and the width of the grooves controls the amount of material removed from the specimen surface. Hardness of the wearing surface with respect to the abrasives and the load exerted controls the extent of the above listed parameters[63-64]. With the increase in hardness of the composite due to aging as well as the formation of the hard precipitates, the aged composite is expected to provide resistance against this action. When compared to the alloy, these

75 reinforcements would carry most of the load and the penetration of the abrasive would be restricted by these hard reinforcements. Microcracking occurs when there is an increase in the stress concentration in selective areas of the wearing surface. Internal notches in the form of pores, microcracks, inclusions etc are the cause of increase in the stress concentration. In this case, large wear debris are detached from the wearing surface due to crack initiation and propagation. According to the wear model described by Zum

Ghar[109], the extent of cracking depends on the loading conditions, the size and the shape of the penetrating abrasive particles, the density and the type of internal notches and the fracture toughness of the wearing material. The increase in hardness due to aging would lead to a decrease in the fracture toughness of the aged composites [110-111]. In addition, the state of stress may be locally intensified by the presence of residual stresses arising from the mismatch of thermal expansion coefficients of the matrix and the reinforcement during quenching. The residual stresses would superimpose with the applied stresses, producing heightened stress concentration at the defects. Keeping in mind the above facts, the probability of microcracking increases in the aged composites.

But in order to predict the true wear rate we also have to consider the mechanism of particle removal.

5.3.3 Mechanism of Particle Removal

The ease of particle removal is expected to be the rate determining step in this study since they protect the matrix from the cutting action of the abrasive. As it was seen from figure 4.1a in the previous section that the composite had a good amount of silicon carbide reinforcements and exhibited a good interfacial bonding. Note that the

Fig. 5.2. SEM Micrograph showing the fracture of the particle

Figure. 5.3. SEM Micrograph of the composite showing weakening of the interfacial bond between the matrix and the particle

Figure. 5.4. SEM Micrograph of the composite showing a pull-out of a SiC particle from the matrix

78 micrograph shown in figure 4.1 was grinded with 800 grit SiC abrasive before being polished. However, examination of the surface after wear suggests a weakening in the interfacial bonding and a decrease in the number of SiC particles on the surface as can be seen from figures 5.3 and 4.11 respectively. This finding suggests that SiC particles have been pulled out of the wear surface in large number and are not offering much resistance to wear. In the present study since the size of the SiC abrasive is quite large ( 268µm) as compared to the reinforcement size (5-20µm) the depth of penetration is expected to be quite high which will reduce the efficiency of the reinforcement particles in protecting the matrix a phenomenon discussed earlier. In addition, the continuous cutting action of the abrasive would remove the matrix from the adjacent areas of the reinforcement leading to the weakening of the interfacial bond and eventually resulting in the pull-out of the SiC particle as shown in figure. Moreover, because of the lattice straining in the surrounding areas of the particles, there will be a reduction in the extent of plastic deformation that these areas can undergo, which will make them more susceptible to cracking. These cracks will result in the removal of the matrix from adjacent areas of the particles, thereby decreasing the strength of interfacial bond. Some of the particles also underwent fracture as can be seen from the figure 5.2. These fractured particles must also become detached from the matrix. In both the above cases, the strength of the bond between the matrix and the particle is expected to play a critical role in determining wear.

Particle

Precipitate PrecipitatePrecipPrecipitateitate

Particle

Precipitate

Fig.5.5. SEM Micrograph of the Aged Composites showing (a) the adheshion of the

precipitate to the particle at a lower magnification (b) adheshion of the precipitate to the

particle at a higher magnification

5.3.4 Dominant Mechanism

From the above discussion it is clear that the removal of the particles from the matrix will play a major role in enhancing wear. In this study, precipitates were observed to be attached to the SiC particles in the aged composites as shown in figure 5.5. The reason for formation of precipitates near the interface can be attributed to the high density of dislocations near the interface. The difference in the thermal expansion coefficient Δα between aluminum and SiC is of the order of 4.7. When the composite is cooled from elevated temperature (during solidification or solution treatment), misfit strains occur due to differential thermal contraction at the matrix/reinforcement interface that are sufficient to generate dislocations. These dislocations can serve as heterogeneous sites for the nucleation of precipitates and can provide short circuit diffusion paths for solute atoms[55]. In addition, the interface between the SiC particle and aluminum matrix would itself act as a good heterogeneous nucleation site for the formation of precipitates.

The adhesion between the interface and precipitates is expected to provide resistance to the SiC particles from scooping off from the composite during wear and hence providing an increase in the wear resistance. This mechanism is more clearly depicted in figure 5.6.

The precipitates which are seen to be attached to the particle will have a certain degree of bonding with the matrix. These attached precipitates will increase the interfacial area in contact with the matrix. Higher the interfacial area more will be the energy required in extracting the particles. Secondly, as seen from figure 5.2, some of the particles got fractured as the pressure increased. The adhering of the precipitates to the matrix will increase the surface area of the particles which would in turn cause a decrease in the

PARTICLE

PRECIPITATE

MATRIX

Figure.5.6. Schemetic showing adherence of the precipitates with the reinforcement

pressure acting on the particles and thereby helping in reducing the rate of crack initiation and propagation to a certain extent. The cracked particles will be more easily removed and would not provide much support to the adjacent matrix, thereby increasing wear.

Thirdly and most importantly, as can be seen from figure 5.6 these adhering precipitates will increase the particle dimensions. As discussed earlier the ratio of the depth of cuts to the reinforcement size would have an effect on wear behavior. When the depth of the cuts is greater than the vertical dimension of the silicon carbide reinforcement, then the reinforcement will not offer any resistance and would be removed along with the matrix.

This mechanism is expected to be a major contributor to wear since the abrasive size is quite large (268µm) as compared to the reinforcement size (5-20µm) and after the wear tests not much SiC reinforcements were seen on the surface implying that the depth of penetration of the abrasive was greater than the reinforcement size in most cases. This mechanism can be more easily understood from figure 5.7. In the figure two particles can be seen, one in which no precipitates are attached and the second one is seen to have adhered with the precipitates. The depth of penetration of the abrasive is little greater than the particle size. As the abrasive approaches the particles, the matrix being softer would not offer much resistance and would be removed by the cutting action of the abrasive. Since the depth of penetration is greater than the particle size, the matrix underneath the particle would be removed and in the process take away the particle along with it. But with the adherence of the precipitates, the size of the particle increases, considering the attached precipitate and the particle as one system. Now as the abrasive

Abrasive Particle Reinforcement particle

Precipitate

Figure 5.7 Schematic showing the interaction of the abrasive with the SiC particle

84 moves towards this particle, the precipitates being harder than the matrix would resist the cutting action to a greater extent and the particle would remain intact in the matrix.

The importance of the ratio of penetration depth and reinforcement size was also discussed by Wang and Rack [108]. Figure 5.9 shows a plot of relative wear resistance of the composite reinforced with SiCp against relative penetration depth hr, where hr is the ratio of average abrasive particle penetration depth (h) and reinforcement diameter (d) as can be seen from the figure 5.8. It shows that above the ratio of unity wear resistance decreases significantly and particles do not offer much resistance and will be easily removed. Since a large number of precipitates were observed in the composite aged which was aged for 5 hours (peak aged), it is highly probable that it will have maximum number of precipitates attached to the particles. Infact as can be seen from figure 5.4b the the density of the precipitates is so much that in many areas they form a network connecting two particles. This was followed by the composites aged for 12 and 40 hours.

No precipitates were observed in the as cast composite and hence it showed maximum wear loss.

Figure 5.8. Schematic diagrams showing interactions between an abrasive particle and a

SiC, composite at penetration depths of (a) h < d and (b) h> d. [108]

Figure 5.9. The relative wear resistance vs. relative abrasive penetration depth.[108]

5.3.5 Comparison of wear behavior of Composite with Alloy

In the present study, it was also observed that on all loads the wear loss in composites was less as compared to the matrix alloy. This finding suggests that even though there has been a good amount of reinforcement pull-out, but they are still present in sufficient number to provide resistance from the cutting action of the abrasives and protect the matrix. Comparison of the worn surfaces of the matrix of the composites with that of the alloy as shown in figures 4.12 and 4.13 suggests that there has been a good amount of plastic deformation in alloy, whereas this deformation is subdued in composites. In the case of Al-Si alloy, it is considered that SiC abrasives could penetrate easily into the soft matrix alloy, resulting in extensive plastic deformation and excessive material removal from the worn surface. In the composites, although very few, but some particles were seen to be protruded on to the surface after the soft matrix was removed due to abrasion as can be seen from figure: 5.10. The hard SiC particles resist the penetration and microcutting action of abrasives effectively, thus reducing the wear loss.

Also plastic deformation was not observed in case of composites. It was also found that heat treatment of alloys did improve the wear behavior of alloys as compared to the as cast alloy, but as compared to the composite the wear rate was still high.

5.3.6 Comparison with Previous Studies

Although much work has been done in the field of abrasive wear behavior of metal matrix composites, but only a limited research has been done on the effect of ageing on wear resistance. In the study of age hardenable Al-Zn-Mg composite containing SiC particulate Lin and Liu [18] found that that the wear resistance was higher in the over-aged condition due to the formation of precipitates over and around the SiC

Fig. 5.10. SEM Micrograph showing a protruding SiC particle

reinforcement. This was believed to have caused an increase in interfacial strength between the particulate reinforcement and the matrix which resulted in improved wear resistance. However, Wang and Rack [112] found a different response. In their study of the effect aging on the wear behavior of SiCw/2124 it was found that the wear resistance decreases when going from the peak aged to the over-aged condition due to the change in matrix microstructure and reduction in hardness. Song[19] also found the similar result.

He attributed this to the coarsening of precipitates which caused the decrease in hardness and wear resistance in 2014/ Al/SiC and 6061 Al/SiC composites. He also found that wear resistance of under-aged composites was less than peak aged which was due to the inability of the precipitates formed during the early stages in providing resistance to plastic deformation as the abrasive particles pass across the composite surface. In the present study, the result obtained was similar to that of Wang and Rack[112] and

Song[19]. However, in those studies the role of adhesion of the precipitates to the interface was not considered. In the present study, the decrease in wear resistance of the overaged composites as compared to the peak-aged one cannot be attributed to the decrease in hardness since cracking is occurring. The increase in hardness of the peak- aged composite will reduce its fracture toughness and make it more susceptible to cracking. The adhesion of the precipitates with the interface is expected to play a critical role in wear behavior.

6.0 Conclusions

1) The hardness increased initially with increasing time. It reached a maximum at

around 5 hours. Further aging reduced the composite hardness.

2) The microstructures of the composites obtained after aging were in agreement

with the aging curve.

3) The abrasive wear resistance of Al/SiCp composite was much higher than that of

the unreinforced Al-Si (10wt %) alloy.

4) The abrasive wear resistance of Al-Si(10wt%) alloy was found to increase with

aging, but no variation in wear rate was found among the aged alloys.

5) The abrasive wear resistance of aged alloys was found less as compared to the as

cast and aged composites

6) The abrasive wear of both the composite and the alloy increases with increasing

load.

7) The abrasive wear resistance of the composites can be altered by aging with the

composites aged to peak hardness exhibiting the maximum wear resistance. Wear

resistance was found to decrease with overaging.

8) Than main mechanisms of matrix removal in the composite were found to be

microcutting and microcracking.

9) Particle removal mainly depended upon the ratio the depth of the groove and

vertical dimension of the reinforcement.

10) Weakening of the interfacial bond between the matrix and the reinforcement due

cutting and cracking of the matrix in the adjacent areas of the reinforcement also

caused particle removal.

11) Fracture of the SiC reinforcement was also found to occur.

12) Microcutting, microcracking along with plastic deformation were the main

mechanisms of wear operating in the unreinforced Al-Si(10wt%) alloy

13) Some of the SiC particles got protruded onto the surface resisting the cutting

action of the abrasive.

14) The dominant mechanism which resulted in the improvement of wear resistance

with aging was found to be the adhesion of the of the Si precipitates with the

interface of the matrix and the reinforcement.

7.0 Future Work

1) Wear tests can be carried out using finer abrasive grits as counterface to study the

contribution of the precipitates in resisting the cutting action.

2) A study of the change in fracture toughness with age hardening needs to be carried

out to study the effect of microcracking in wear behavior.

3) A wear study of the under-aged and the over-aged composites having the same

hardness can be carried out to study the role of hardness and size of precipitates.

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